Flow Induced Crystallization of Polyolefins

Flow induced crystallization of polyolefins PROEFSCHRIFT ter verkrijging van de graad van doctor aan de Technische Universiteit Eindhoven, op gezag van de Rector Magnificus, prof.dr.ir. C.J. van Duijn, voor een commissie aangewezen door het College voor Promoties in het openbaar te verdedigen op woensdag 16 januari 2008 om 14.00 uur door Luigi Balzano geboren te Pompeï, Italië Dit proefschrift is goedgekeurd door de promotoren: prof.dr. S. Rastogi en prof.dr. P.J. Lemstra Copromotor: dr.ir. G.W.M. Peters A catalogue record is available from the Eindhoven University of Technology Library ISBN: 978-90-386-1199-0 Copyright © 2008 by L. Balzano Printed at the Universiteitsdrukkerij, Eindhoven University of Technology, Eindhoven. Cover design: Luigi Balzano and Bregje Schoffelen (Oranje Vormgevers) The research described in this dissertation was financially supported by the Dutch Polymer Institute. (DPI) project # 132. Table of Contents  Summary   ......................................................................................................................... 1  Chapter 1 Introduction ...................................................................................................... 3  1.1  Preamble ..................................................................................................... 3  1.1.1  Process‐properties relation ............................................................ 3  1.1.2  Historical survey on polyethylene and polypropylene ................... 4  1.2  Processing of polyolefins  ............................................................................ 5  . 1.3  Aim of the thesis ......................................................................................... 7  1.4  Outline of the thesis .................................................................................... 8  1.4.1  iPP melts containing small amount of DMDBS .............................. 8  . 1.4.2  PE melts with a bimodal molecular weight distribution ................ 9  1.4.3  iPP with unimodal molecular weight distribution .......................... 9  1.5  References ................................................................................................. 10  Chapter 2  Flow induced crystallization in iPP‐DMDBS blends: implications on morphology  of shear and phase separation ....................................................................... 13  2.1  Introduction .............................................................................................. 14  2.2  Experimental method ............................................................................... 16  2.2.1  Materials ....................................................................................... 16  2.2.2  Sample preparation ...................................................................... 17  2.2.3  X‐Ray characterization .................................................................. 17  2.2.4  Rheological characterization ........................................................ 19  2.2.5  DSC  ............................................................................................... 19  . 2.3  Results and discussion .............................................................................. 20  2.3.1  Effects of DMDBS on structure and morphology of iPP in the solid  state .............................................................................................. 20  2.3.2  Crystallization under quiescent conditions .................................. 22  2.3.3  Morphology of the system in Region II ........................................ 26  i 2.3.4  Rheology of the system in Region II ............................................. 29  2.3.5  Effect of flow on iPP‐DMDBS blends near the gel transition ....... 31  2.3.6  Morphological implications of flow and DMDBS phase separation  on the crystallization of iPP .......................................................... 34  2.4  Conclusions ............................................................................................... 36  2.5  References ................................................................................................. 38  Chapter 3  Thermo‐reversible DMDBS phase separation in iPP: effects on flow induced  crystallization ................................................................................................. 41  3.1  Introduction .............................................................................................. 41  3.2  Experimental method ............................................................................... 43  3.2.1  Materials ....................................................................................... 43  3.2.2  Sample Preparation ...................................................................... 43  3.2.3  X‐Ray Characterization ................................................................. 44  3.2.4  Rheological Characterization ........................................................ 45  3.2.5  DSC  ............................................................................................... 45  . 3.3  Results and discussion .............................................................................. 45  3.3.1  Thermoreversibility in the phase diagram ................................... 45  3.3.2  Linear viscoelasticity of the system in Region I‐PS ....................... 52  3.3.3  Crystallization on cooling after flow in Region I‐PS  ..................... 53  . 3.4  Conclusions ............................................................................................... 58  3.5  References ................................................................................................. 59  Chapter 4   Crystallization and dissolution of flow induced precursors ............................ 63  4.1  Introduction .............................................................................................. 63  4.2  Experimental method ............................................................................... 65  4.2.1  Synthesis of a bimodal HDPE ........................................................ 65  4.2.2  X‐ray characterization .................................................................. 66  4.3  Results and discussion .............................................................................. 67  4.3.1  Rheological characterization ........................................................ 67  ii 4.3.2  Thermodynamics of flow induced precursors .............................. 70  4.3.3  Flow induced precursors just above the equilibrium melting  temperature ................................................................................. 72  4.4  Conclusions ............................................................................................... 79  4.5  References ................................................................................................. 79  Chapter 5  Precursors, crystallization and melting in sheared bimodal HDPE melts ......... 83  5.1  Introduction .............................................................................................. 84  5.2  Experimental method ............................................................................... 85  5.2.1  Material preparation .................................................................... 85  5.3  Characterization........................................................................................ 85  5.3.1  Rheology ....................................................................................... 85  5.3.2  Small Angle X‐Ray Scattering (SAXS). ........................................... 86  5.3.3  Wide Angle X‐Ray Scattering (WAXS or WAXD). .......................... 86  5.3.4  Shear experiments  ....................................................................... 87  . 5.4  Results and discussion .............................................................................. 87  5.4.1  Flow induced precursors above the equilibrium melting  temperature ................................................................................. 87  5.4.2  Stable and relaxing precursors above the equilibrium melting  temperature ................................................................................. 90  5.4.3  Flow induced shishes below the equilibrium melting  temperature: the influence of temperature ................................ 91  5.4.4  Flow induced shishes below the equilibrium melting  temperature: the influence of flow conditions ............................ 93  5.4.5  Separating shish creation from the kebab crystallization ............ 95  5.4.6  Crystallization onset temperature after short term shear ......... 100  5.4.7  Melting of shish kebabs .............................................................. 101  5.5  Conclusions ............................................................................................. 103  5.6  References ............................................................................................... 104  iii Chapter 6  Metastable structures during fast short term shear ..................................... 107  6.1  Introduction ............................................................................................ 107  6.2  Materials and methods ........................................................................... 109  6.2.1  Materials ..................................................................................... 109  6.2.2  X‐ray characterization ................................................................ 109  6.2.3  Shear experiments  ..................................................................... 111  . 6.3  Results and Discussion ............................................................................ 112  6.3.1  Flow conditions in short term shear .......................................... 112  6.3.2  Flow induced precursors during short term shear ..................... 113  6.3.3  Crystallization after short term shear. ....................................... 114  6.4  Conclusions ............................................................................................. 121  6.5  References ............................................................................................... 121  Chapter 7 Conclusions and recommendations ............................................................... 125  7.1  Conclusions / Technology assessment .................................................... 125  7.2  Recommendations for future research ................................................... 126  Samenvatting ................................................................................................................ 129  Acknowledgements ....................................................................................................... 133  Curriculum Vitae ........................................................................................................... 135        iv Summary  Flow induced crystallization of polyolefins  Polymers are a widespread class of materials that provide an often advantageous combination of properties. Easy processability and high versatility combined with low costs make polymers the materials for an increasing number of high-tech and commodity applications. Semi-crystalline polyolefins are an important class of polymers, produced in more than 150 million metric tons per year. They are used to make a wide range of products ranging from fibers with superior mechanical properties to flexible packaging and molded parts. The properties of these materials are related to the whole history of the material, from chemistry/catalysis in the reactor and, in particular, to the processing conditions. Nowadays, there is a growing interest in added value to these products by achieving outstanding properties such as high stiffness (up to ~150 GPa) for fibers and clarity for injection molded parts. This demands more thorough studies on the process-properties relation that is not yet fully understood. The main objective of this thesis is to enhance the nucleating efficiency of the polymer by inducing oriented structures possessing good epitaxial matching. The goal is achieved by developing the oriented structures in polymer melt either by (a) making use of fillers that self assemble into nano sized fibrils and orient under flow or (b) by the addition of identical higher molar mass molecules possessing considerably higher relaxation times compared to the base (matrix) polymer. In the first part of the thesis, the crystallization of isotactic polypropylene (iPP) in the presence of 1,3:2,4-bis(3,4-dimethylbenzylidene)sorbitol (DMDBS) is discussed. DMDBS is a small organic compound with a high melting temperature (~250 °C) used as a nucleating agent and so-called clarifier for iPP. The nucleating efficiency of this compound, in the low concentration regime (less than 1 wt%), is very high and leads to very small iPP crystallites 1 that confer clarity to the material. DMDBS can crystallize within the molten polymer matrix forming a percolated network of nano-fibrils whose surface hosts a large number of tailored nucleation sites. Because of the epitaxial relation between iPP and DMDBS, iPP lamellae grow always radially on DMDBS fibrils, i.e. with the crystalline c-axis parallel to the fibril axis, the so the so-called shish-kebab morphology (rather similar to the well-known food product). Therefore, the orientation of DMDBS fibrils templates the orientation of iPP lamellae. Randomly oriented DMDBS fibrils lead to randomly oriented iPP lamellae and aligned DMDBS fibrils lead to aligned iPP lamellae. A long lasting alignment of DMDBS fibrils can be obtained deforming their network even above the melting point of the polymer. Nearly 0.5 wt% of oriented DMDBS fibrils can template very oriented (fiber-like) polymer morphologies. In the second part of this thesis, the flow induced crystallization of high density polyethylene (HDPE) with a bimodal molecular weight distribution is discussed. This material is an intimate blend of low and high molecular weight polymer chains (LMW and HMW) synthesized with a new chemistry route as described in the thesis of N. Kukalyekar (Ph.D. thesis Eindhoven University of Technology, December 2007). Just above the equilibrium melting temperature (T 141.2  ) of the polymer, the mutually entangled HMW chains can be stretched with shear and, due to the restricted number of molecular conformations, nucleate into needle-like crystals. By choosing appropriate flow conditions, a suspension of shishes (extended chain crystals) can be formed while the nucleation of kebabs (folded chain crystals) is suppressed because of a too high temperature. With perfect epitaxy matching and a good state of dispersion, shishes are the ideal substrate for the nucleation of HDPE lamellae. On cooling after the application of shear at 142 , HDPE lamellae nucleate using shishes as a heterogeneous substrate and, therefore, with the crystalline c-axis parallel to the shish direction. Similarly to the case of iPP-DMDBS blends, it is observed that nearly 0.5 wt% of pre-aligned shishes can template very oriented (fiber-like) morphologies. 2 Chapter 1  Introduction  1.1 Preamble  1.1.1 Process‐properties relation  Polymeric materials exhibit an intricate process-properties relationship that links the properties of the final products to the whole history of the material. Figure 1.1 describes the connections from synthesis, via processing, to product properties. Figure 1.1: Flow chart describing the process-properties relationship in semi-crystalline polymeric materials. In the last decades, many scientific studies have been devoted to identify relevant parameters and their role in this relationship. Interdisciplinary efforts have led to the production of new materials, with advantageous properties, that have replaced traditional ones (glass, ceramics, metal, wood, …) in many applications and have enabled developments in new areas, like micro-electronics and biomedical applications. However, some aspects of the process-properties relation in polymeric materials are not yet fully understood. Their clarification could lead, eventually, to materials with properties tailored to the application. A 3 Chapter 1 modern shift in industrial paradigms demands to achieve this goal without developing ‘new’ polymers but, instead, making use of ‘old’ polymers that are based on relatively cheap and readily available monomers1. For many applications, polyolefins are the ideal candidates. 1.1.2 Historical survey on polyethylene and polypropylene  Polyolefins are commodity materials obtained by polymerizing olefins (alkenes-1). Nowadays, with more than 150 million metric tons2 per year, polyolefins are the most widespread class of polymers. Polyolefins are inert materials and, when recycled, environmentally harmless. The number of products based on these materials, from packaging to ballistic, from structural to biomedical applications, increases every day. The simplest polyolefin is polyethylene (PE) that is a sequence of ethylene monomers. Polyethylene was discovered in the 1930s by Fawcett and Gibson at ICI in strong collaboration with prof. T. Michels of the Free University of Amsterdam who pioneered the behavior of gases at elevated pressures. The first industrial PE grades were produced by the English company ICI in 1939. Initially, it was possible to produce only a highly branched and with low density PE (LDPE). This highly amorphous material, with high toughness, is still used in today’s packaging applications. A major breakthrough came in the 1950s, when Ziegler3 and Natta4, 5 (1963 Nobel Prize in Chemistry laureates) discovered organometallic catalysts capable of synthesizing high density linear polyethylene (HDPE). Because of a regular chain structure, HDPE can partially crystallize and it exhibits better mechanical properties. In the 1960s, polyethylene attracted the attention of physicists. Pennings6 and Keller7 pioneered the formation of elongated crystals (shish-kebabs) in stirred solution and stressed melts. At the same time, Ward9 found that upon solid state drawing of melt crystallized HDPE, re-organization of the molecules increases the E-modulus up to 60 GPa. These studies unveiled the role of the morphology in the properties of semicrystalline materials, enabling developments in the area of high performance materials from flexible molecules. At the end of the 1970s, at DSM Research in the Netherlands, Smith and Lemstra10, 11 invented a process to spin ultra high molecular weight PE (UHMWPE) from a semi-dilute solution. After drawing, E-moduli of up to 150 GPa could be achieved. One of the last breakthroughs was at the end of the 1970s, when Kaminsky12 discovered metallocene 4 Chapter 1 catalysts allowing for narrow molecular weight distributions and enhancing the control over chain structure of homo- and co-polymers. The second simplest polyolefin is polypropylene (PP) that is obtained by polymerizing propylene. PP is basically PE with a methyl side group every other carbon atom in the back-bone. The relative orientation of the side groups in the space (tacticity) is very important for the properties of the material. Atactic PP (aPP), with randomly distributed side groups, can not crystallize and is a rubbery material. In contrast, isotactic PP (iPP), with the side groups consistently on one side, has the necessary long range order required for crystallization. iPP is a competitor for HDPE because it has a higher melting point and can be made transparent with the use of clarifying agents. The synthesis of iPP was enabled by Ziegler-Natta catalysts13 and was performed, for the first time, by the Italian company Montecatini in 1957. 1.2 Processing of polyolefins  This thesis deals with topics closely related to melt-processing of polyolefins. Polyolefins are often processed via the molten state, applying flows and temperature gradients. Melt-processing has the advantage of not involving solvents and can be used to create complicated shapes. However, with the design of a manufacturing process for polyolefins and, more in general, for all polymeric materials, one should also consider parameters like molecular weight (Mw) and molecular weight distribution (MwD). It is well established that the viscosity of the melt (η scales with Mw according to a power law14-18: M . . Melt-processing is possible only for relatively low molecular weight materials. In the other cases, more complicated routes are available but, often, they are limited to simple profiles, mostly fibres and tapes. Flow during processing enhances the crystallization rate of the polymer by promoting the formation of nuclei of the crystalline phase19-33. This alters the final morphology of the polymer and thus the (mechanical, optical, transport, …) properties of the material34, 35 . Remarkably, the final morphology of the polymer strongly depends on the structures, called precursors, present in the early stages of crystallization. These precursors are structures with undetectable degree of crystallinity but with a certain degree of order. Occasionally, because 5 Chapter 1 of local density fluctuations or further growth due to flow, precursors exceed some critical dimensions and become spontaneous growing crystalline nuclei. Flow induced precursors (FIPs) can be generated at relatively high temperatures (i.e. around the thermodynamic melting point) and, for a strong enough flow, exhibit an anisotropic morphology. These precursors can be quite large and they initiate the growth of ‘shish kebabs’; i.e. anisotropic crystallites made of a fibrillar core decorated with a stack of lamellae 6, 8, 36 . In some cases, shish-kebabs can entirely replace the spherulitic assemblies of lamellae that are characteristic for crystallization in quiescent conditions. This can be advantageous for some polymer products but, definitely, not for all of them. For instance, shish-kebabs cause a high modulus and a high strength in fibres10, weakness (brittleness) 39,40 37, 38 . In contrast, they can be the source of mechanical in injection-moulded products. Figure 1.2 shows examples of spherulites and shish-kebabs obtained by crystallizing polyolefins under quiescent and flow conditions. Figure 1.2: a) Scanning electron micrograph of a melt crystallized iPP spherulites (courtesy P.Schmit); b) Optical micrograph of iPP spherulites growing in the melt (reproduced with permission from Figure 2, page 32 of reference 41); c) polyethylene shish-kebab in the melt (reproduced with permission from reference 42); d) Polyethylene shish-kebabs forming zip fastener structures (reproduced with permission from reference 43); e) Multiple shishes crossing the same kebabs in polyethylene (reproduced with permission from Figure 3 of reference 44). 6 Chapter 1 For some polymers, for instance iPP, nucleation is relatively slow and processing is accelerated with nucleating agents (NAs) 45-47. NAs have a marked impact on the morphology of the polymer and, by reducing the size of the crystallites, can improve the mechanical properties and reduce the haze. When using NAs, their chemical nature, concentration, dispersion and aspect ratio need to be considered as extra parameters affecting the final morphology of the polymer. In addition, during flow, the nucleating particles influence the local distribution of stresses enhancing the orientation in the surrounding molecules. This phenomenon can be very important in the flow induced crystallization of polymer melts containing fillers48. 1.3 Aim of the thesis  The aim of this thesis is to identify basic principles for the onset of oriented morphologies (shish-kebabs) during flow of melt-processable semicrystalline polyolefins. For polyolefins, the objective is often to achieve the desired properties with melt processing at low costs. When the desired properties are the result of an oriented morphology, the goal can be attained by a) tailoring the melt with small amounts of a ‘smart’ additive or b) with a clever choice of the molecular weight distribution, both in combination with the right processing conditions (temperature and flow history). In our experimental work, we consider three systems: • iPP containing small amount of 1,3:2,4-bis(3,4-dimethylbenzylidene)Sorbitol or DMDBS; • PE with a bimodal molecular weight distribution; • iPP with unimodal molecular weight distribution; In all cases, molecular weights allowing melt processability are selected. 7 Chapter 1 1.4 Outline of the thesis  1.4.1 iPP melts containing small amount of DMDBS  DMDBS is an additive which is used as a nucleating agent for iPP 49, 50. The formula is shown in Figure 1.3. The affinity between this additive and the polymer is very high. Only tiny amounts of DMDBS, less than 1 wt%, cause dramatic changes in the morphology of iPP. Crystal assemblies can become smaller than the wavelength of the light (~400 nm) and turn iPP from an opaque to a clear and transparent material 51. Figure 1.3: Chemical structure of DMDBS. The polar molecules of DMDBS can dissolve in the molten iPP only at very high temperatures. On cooling, DMDBS self-assembles, phase separating from the melt, and forms a percolated network of fibrils 52 whose surface hosts nucleation sites tailored for iPP. The state of the art regarding the crystallization of iPP in presence of small amount of DMDBS is described in the Introduction to Chapter 2 and Chapter 3. In Chapter 2*, the impact of DMDBS on the crystallization of iPP is discussed. In particular, we address the role of DMDBS fibrils in templating the iPP morphology after flow (shear) at high temperatures where the viscosity of the melt is low and the relaxation times are short. In Chapter 3†, the thermo-reversibility of DMDBS phase separation is studied and the investigation on the role of DMDBS fibrils in templating the iPP morphology is extended to higher temperatures. * Partially reproduced from: Balzano, L. et al. ‘Flow induced crystallization in iPP-DMDBS blends: implications on morphology of shear and phase separation’, Macromolecules 2007 (Accepted) † Partially reproduced from: Balzano, L. et al. ‘Thermo-reversible DMDBS phase separation in iPP: effects on flow induced crystallization ’, Macromolecules 2008 (Submitted) 8 Chapter 1 1.4.2 PE melts with a bimodal molecular weight distribution  It is well established that small amounts of high molecular weight chains promote the formation of shish-kebabs during flow induced crystallization 22, 26, 28, 30, 33, 53-58 . It has been proposed that the underlying mechanism relies on the enhanced creation of flow induced precursors with anisotropic morphology due to stretching of the high molar mass chain network 59-61 . Stretched chains have a high segmental orientation that allows them to crystallize faster than coiled chains 19, 62 and form fibrillar crystals 21 . To validate this hypothesis, we use a specially synthesized blend of a low molar mass linear HDPE containing 7 wt% of high molar mass linear HDPE. Under shear, this material exhibits a high tendency to generate shish-kebabs. Shish generated at high temperature can be used as a heterogeneous substrate for the nucleation of the rest of the molecules; exact lattice matching and good state of dispersion make them the ideal nucleating substrate. The state of the art of flow induced crystallization, relevant to the work presented in this thesis, is summarized in the Introduction paragraphs of Chapter 4 and Chapter 5. In Chapter 4*, the dynamics of flow induced precursors just above the equilibrium melting point is discussed. This investigation unveils, for the first time, the possibility to generate a suspension of extended chain shishes only. In Chapter 5†, the investigation on the nature of shishes just above the equilibrium melting temperature is expanded and their potential as nucleators for the bulk of the polymer is systematically explored. 1.4.3 iPP with unimodal molecular weight distribution  In Chapter 4 and 5, the formation of shishes via crystallization of needle-like flow induced precursors is observed at relatively high temperature after the application of shear. 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Polymer 2004, 45, 3249-3256. 12 Chapter 2*  Flow induced crystallization in iPP‐DMDBS  blends: implications on morphology of shear  and phase separation  Nucleation is the limiting stage in the kinetics of polymer crystallization. In many applications of polymer processing, nucleation is enhanced with the addition of nucleating agents. 1,3:2,4-bis(3,4-dimethylbenzylidene) sorbitol or DMDBS is a nucleating agent tailored for isotactic polypropylene (iPP). The presence of DMDBS changes the phase behavior of the polymer. For high enough temperatures the system iPP-DMDBS forms a homogeneous solution. However, in the range of concentration spanning from 0 to 1 wt% of DMDBS, the additive can phase separate/crystallize above the crystallization temperature of the polymer, forming a percolated network of fibrils. The surface of these fibrils hosts a large number of sites tailored for the nucleation of iPP. The aim of this Chapter is to investigate the combined effect of flow and DMDBS phase separation on the morphology of iPP. To this end, we studied the rheology of phase separated iPP-DMDBS systems and its morphology with time resolved Small Angle X-ray Scattering (SAXS). The effect of flow is studied combining rheology, SAXS and a short term shear protocol. We found that, with phase separation, DMDBS forms fibrils whose radius (~5 nm) does not depend on the DMDBS concentration. The growth of these fibrils leads to a percolated network with a mesh size depending on DMDBS concentration. Compared to the polymer, the relaxation time of the network is quite long. A shear flow, of 60 s-1 for 3 s, is sufficient to deform the network and to produce a long-lasting alignment of the fibrils. By design, lateral growth of iPP lamellae occurs orthogonally to the fibril axis. Therefore, with crystallization, the preorientation of DMDBS fibrils is transformed into orientation of the lamellae. This peculiarity is used here to design thermo-mechanical histories for obtaining highly oriented iPP morphologies after shearing well above the melting point of the polymer (i.e. without any undercooling). In contrast, when shear flow is applied prior to DMDBS crystallization, SAXS showed that iPP crystallization occurs with isotropic morphologies. * Partially reproduced from: Balzano, L. et al. ‘Flow induced crystallization in iPP-DMDBS blends: implications on morphology of shear and phase separation’, Macromolecules 2007 (Accepted) 13 Chapter 2 2.1 Introduction  Morphology control is an important issue in polymer processing as it influences a broad range of properties of the final products. For instance, mechanical, optical and transport properties of polymeric materials depend on the size and shape of the crystallites1, 2. It is well known that thermal and mechanical histories do play an important role in the creation of these morphological features3, 4 and that additives can also have a remarkable influence2, 5-8. Nucleating agents are a family of additives used to speed up processing rates of polymers. In the case of isotactic polypropylene (iPP) a common nucleating agent is a sorbitol derivative: 1,3:2,4-bis(3,4-dimethylbenzylidene)sorbitol or DMDBS. DMDBS is a chiral molecule that can self-assemble or crystallize within the molten polymer matrix. Self-assembly takes place because of inter-molecular hydrogen bonds formation. Hydrogen bonds, in this case, work essentially in one direction and drive the molecules to pile up (see Figure 2.1). This leads to the unidirectional growth of fibrillar crystals. Elementary DMDBS fibrils, in iPP, have a diameter of ~10 nm and a length up to several microns. They can also form bundles with a diameter up to 100 nm. Figure 2.1: A stack of two DMDBS molecules. Crystallization of DMDBS within the iPP matrix corresponds to a liquid-solid phase separation, in the following, referred to as DMDBS crystallization or DMDBS phase separation. The DMDBS molecule has a special ‘butterfly’ configuration, see Figure 1.3. The ‘wings’ of the molecule (phenyl rings with two methyl groups attached) enable dissolution in the polymer and, at the same time, are tailored nucleation sites for iPP, while the ‘body’ 14 Chapter 2 comprises two moieties: one dictates the geometry of the molecule and the other bears the polar groups (hydroxyls) for hydrogen bond formation9. Polarity is one of the main features of DMDBS. In contrast, iPP is a fully apolar molecule. This difference becomes clear and leads to a rich phase diagram when iPP and DMDBS are compounded together. Kristiansen et al.10 proposed a monotectic model for this phase diagram where the eutectic point lies near 0.1 wt% of the additive. In their model, miscibility of the two molecules is always possible at high temperatures (Region I). They define four concentration regimes based on different phase transitions occurring during the cooling of a homogenous mixture. From the application point of view, the most interesting concentration regime extends from ~0.1 wt% to ~1 wt% of DMDBS where iPP exhibits a high clarity. The phase diagram, in this concentration range, is schematically shown in Figure 2.2. Figure 2.2: Schematic phase diagram for iPP-DMDBS mixtures up to ~1 wt% (quantitative data shown in Figure 2.9). Cooling a homogeneous mixture (Region I) leads to crystallization of DMDBS before crystallization of the polymer (Region II) 10. With crystallization, DMDBS forms a percolated network of fibrils suspended in the polymer matrix. The nucleation sites for the polymer reside on the surface of this network. The fibrillar arrangement provides a high surface to volume (S/V) ratio and, therefore, a large number of nucleation sites per unit of volume. However, S/V alone cannot explain the nucleation ability of DMDBS. Thierry et al.9, 11 and Fillon et al.11 demonstrated that DMDBS is a good nucleating agent for iPP because of a good lattice matching between its crystals and the 31 helix of the polymer. The same authors 15 Chapter 2 also define an efficiency scale for nucleating agents, ranging from 0 to 100 %, based on characteristic crystallization temperatures. Dibenzylidene sorbitol (DBS), a nucleating agent very similar to DMDBS, was rated at 41 %. Among several nucleating agents, they found that 4-Biphenyl carboxylic acid (2 wt% in iPP) has the highest nucleation efficiency (66%). The effect of several sorbitol based nucleating agents on quiescent crystallization kinetics and the morphology of iPP has been widely explored12-14, as was the rheology of these systems12, 15, 16. Surprisingly, little attention has been paid to the role of sorbitol based nucleating agents on the crystallization of iPP during or after imposition of a flow, the most common scenario in processing. A notable exception is the work of Nogales et al.17, 18. They studied the flow induced crystallization of iPP-DBS compounds after the phase separation of the additive under well defined conditions, by means of both scattering and imaging techniques. For 1 wt% of DBS, they observed, during cooling, after application of modest shear flows (shear rates ranging from 0.1 to 20 s-1 at 170 ºC), the formation of polymer morphologies characterized by high degrees of orientation. However, the role of DMDBS phase separation in flow induced crystallization of iPPDMDBS blends is not yet fully clarified and is the topic of this Chapter. The work includes also the changes in the rheology of the melt, associated to the formation of the DMDBS fibrillar network, and the flow behavior of this network. The results are based on a combination of Small Angle X-ray Scattering (SAXS), Dynamic Scanning Calorimetry (DSC) and Rheology. Four different iPP-DMDBS blends, containing 0, 0.3, 0.7 and 1.0 wt% of the additive are investigated in quiescent and flow conditions. We address three aspects of these blends: 1. crystallization without application of flow (quiescent conditions); 2. influence of flow prior to the crystallization of DMDBS; 3. influence of flow after crystallization of DMDBS. 2.2 Experimental method  2.2.1 Materials  The iPP used in this work is a commercial homopolymer grade from Borealis GmbH (Austria), labeled HD120MO, with molecular weight, Mw, of 365 kg/mol and a 16 Chapter 2 polydispersity, Mw/Mn, of 5.4. DMDBS (Millad 3988) was obtained in powder form from Milliken Chemicals (Gent, Belgium) and used as received. 2.2.2 Sample preparation  The polymer, available in pellets, was first cryo-ground and then compounded with DMDBS in a co-rotating twin screw mini-extruder (DSM, Geleen) for 10 min at temperatures ranging from 230 to 250 ºC, the higher the DMDBS concentration the higher the compounding temperature used. To prevent degradation of both, polymer and additive, this operation was performed in a nitrogen rich atmosphere. The material obtained was compression molded with a hot press into films of different thicknesses: 1mm for rheology and 200μm for X-ray experiments. The compression molding temperature was 220 ºC and the molding time was 3min. The resulting films were quenched to room temperature and cut in disk-like samples. Following the same procedure, three blends of iPP with 0.3, 0.7 and 1 wt% of DMDBS were prepared. For convenience, these three blends are respectively referred to as B03, B07 and B1 in the text. 2.2.3 X‐Ray characterization  X-ray characterization was performed at the European Synchrotron Radiation Facility (ESRF) in Grenoble (France). Time resolved Small Angle X-ray Scattering (SAXS) experiments were performed at beamline BM26/DUBBLE. Scattering patterns were recorded on a two dimensional gas filled detector (512x512 pixels) placed at approximately 7.1 m from the sample. Scattering and absorption from air were minimized by a vacuum chamber placed between sample and detector. The wavelength adopted was λ=1.03 Å. SAXS images were acquired with an exposure of 5 s and were corrected for the intensity of the primary beam, absorption and sample thickness. The scattered intensity was integrated and plotted against the scattering vector, q = (4π / λ)sin(ϑ / 2) where ϑ is half of the scattering angle. The long period was calculated as Lp = 2π / (qI MAX ) , where qI MAX is the q value corresponding to the maximum in the scattered intensity. Finally, we defined an integrated intensity as: q max II = ∫ I (q )dq where q min and q max are the minimum and the maximum experimentally q min 17 Chapter 2 accessible q values respectively. Two dimensional SAXS images were also used for the characterization of anisotropic morphologies. For this purpose, it was necessary to define three azimuthal regions19. The definitions adopted in the present work are given in Figure 2.3. Figure 2.3: Anisotropic two dimensional SAXS image with definitions of the azimuthal intensity regions. Arrow indicates the applied flow direction. Shear flow experiments in combination with SAXS were carried out in a Linkam Shear Cell (CSS-450) modified with Kapton windows using a ‘short term shearing’ protocol. First, samples were annealed at 230 ºC for 3 min to erase the memory of any previous thermo-mechanical treatment. Next, the temperature was decreased by 10 ºC/min to the desired test temperature where flow was applied under isothermal conditions. For the purpose of this Chapter, we limit ourselves to the application of only one shear condition: nominal shear rate of 60 s-1 for 3 s. Finally, depending on the experimental requirements, the temperature was either decreased to the room temperature or kept constant. Wide Angle X-ray scattering (WAXD) experiments were performed separately on beamline ID11 of the ESRF. The results were used to determine crystallinity and the phases present in the samples. Two dimensional images were recorded on a Frelon detector. Before analysis, the scattering of air and of the empty sample holder was subtracted. After radial integration, the intensity was plotted as a function of the scattering angle 2ϑ . Deconvolution of the amorphous and crystalline scattered intensities was performed using a sixth order polynomial to capture the ‘amorphous halo20, 21. The crystallinity index, a measure of the crystal volume fraction, was calculated as: XWAXD = AC AC + AA 18 ⋅ 100 (2.1) Chapter 2 where, AA and AC are the scattered intensities from the amorphous and the crystalline phases, respectively. 2.2.4 Rheological characterization  Rheological measurements were performed in the linear viscoelastic regime using a strain-controlled ARES rheometer equipped with a 2KFRT force rebalance transducer. In all cases a plate-plate geometry with a diameter of 8 mm was used. Appropriate values of strain were determined with amplitude sweep tests carried out at 5 rad/s over a broad range of strains (ranging from 0.01 to 100 %)22. During the study of phase transitions, large strains can enhance the process and/or affect the morphology23. These effects are minimized by using strains as low as 0.5 % in the experiments. 2.2.5 DSC  The crystallization behavior of the three binary blends iPP-DMDBS was studied in quiescent conditions using Dynamic Scanning Calorimetry (DSC). Samples of approximately 2 mg were placed into aluminum pans and tested in nitrogen atmosphere in a Q1000 calorimeter (TA Instruments). The first step in the thermal treatment was always annealing at 230 ºC for 3 min to erase earlier thermo-mechanical histories. Next, samples were cooled to room temperature at a constant cooling rate of 10 ºC/min. Before identifying peak positions and determining crystallinity, a linear baseline was subtracted from the measured heat flow as a function of the temperature. Finally, crystallinity Te could be estimated as: X DSC 0 c = ΔHc ΔH , where ΔH c = ∫ (dH dT )dT and ΔH c0 are, Ts respectively, the enthalpy of crystallization of the sample and the enthalpy of crystallization of an ideal 100 % crystalline iPP sample (207.1 J g-1) 24. 19 Chapter 2 2.3 Results and discussion  2.3.1 Effects of DMDBS on structure and morphology of iPP in the solid state  In semi-crystalline polymers, structure and morphology depend on the crystallization conditions (thermal and mechanical histories). In order to isolate the effects due to the presence of DMDBS in the solid state, samples were prepared under the same crystallization conditions i.e. quiescent crystallization with 10 ºC/min. Figure 2.4 reports WAXD integrated intensities at room temperature for the neat iPP and the blends with DMDBS. Figure 2.4: WAXD profiles of iPP at room temperature as function of DMDBS concentration. All samples were prepared in the same conditions, i.e. crystallization from the melt at 10 ºC/min. Presence of DMDBS induces the broad 117 peak, indicated by the arrow, that is associated to the formation of γ phase crystals. The crystallinity index is ~60 % in all cases while the amount of γ phase decreases with DMDBS concentration. Note that, curves are shifted in the vertical direction for clarity. The neat iPP shows the typical diffraction peaks of the α crystalline modification. When the additive is present, although the α form remains prevalent, the crystal structure of the polymer shows some specific changes. The 111 peak becomes better resolved and a broad 117 reflection appears. This indicates the simultaneous formation of less defected α and small γ crystals. However, we do not observe significant variation in the WAXD crystallinity index; in all cases, it lies around 60 %. According to Foresta et al.24, the formation of γ phase crystals in presence of the nucleating agent can be explained from a thermodynamic point of view. The nucleating agent shifts the crystallization of the polymer at higher temperatures 20 Chapter 2 where nucleation of γ phase is favored and can compete with nucleation of α phase. The ratio between γ and α phase crystals, X γ , can be estimated with: Xγ = A117 A130 + A117 (2.2) where A130 and A117 are the areas of the non overlapping parts of the 117 and 130 peaks. These two peaks were selected because they are the diagnostic reflections of the γ and the α phase respectively. In the investigated range of concentration, the γ phase content, X γ , is maximum for B03 ( X γ =0.15) and drops for B07 ( X γ =0.09) and B1 ( X γ =0.08). This drop is probably related to a faster α nucleation rate at higher DMDBS concentrations. On the morphological side, the long period of iPP lamellae shows pronounced changes as a function of DMDBS concentration going from 19 nm of the neat sample to 23 nm (average value) of samples containing DMDBS, see Figure 2.5. Figure 2.5: Long periods of iPP lamellae at room temperature as a function of DMDBS concentration. All samples were prepared under the same conditions, i.e. crystallization from melt at 10 ºC/min. The neat iPP shows a long period of 19 nm and this value rises to ~23 nm for samples containing DMDBS. This increase in long period is due to the formation of thicker crystals in presence of DMDBS The lamellar thickness, TL, can be expressed as TL = Lp ⋅ x . Since, the crystallinity index does not vary, our experimental observations are consistent with the formation of thicker crystals when DMDBS is present. The reason for this increase in crystal thickness is 21 Chapter 2 the higher crystallization temperature in presence of the nucleating agent8 that is discussed hereafter. 2.3.2 Crystallization under quiescent conditions  When cooling a homogeneous mixture of iPP and DMDBS to room temperature, two phase transitions are observed: crystallization of DMDBS and crystallization of the polymer. DSC experiments reveal the temperatures and enthalpies characterizing both these transitions. In the cooling thermograms of Figure 2.6 the crystallization peaks of the polymer are, in all cases, clearly visible. A closer look discloses another, much smaller, exotherm at higher temperatures. Figure 2.6: DSC cooling thermograms (after subtraction of a linear baseline) for the neat polymer and blends B03, B07 and B1. Experiments were performed at 10ºC/min, in N2 atmosphere, after annealing the samples at 250 ºC for 3 min. Curves are shifted along the vertical axis for clarity. With the addition of only 0.3 wt% of DMDBS, the crystallization peak shifts to 132 ºC. and its position does not change with further addition of the additive. Nevertheless, the crystallization peak becomes narrower when increasing the amount of DMDBS. This smaller exotherm is associated with the crystallization of DMDBS and, due to the small amount of the additive, becomes visible only after sufficient magnification, see Figure 2.7. Some relevant DSC data during the cooling experiments are summarized in Table 2.1. Note that these data provide enough information to sketch the phase diagram of the system in the investigated range of concentration. 22 Chapter 2 Figure 2.7: Magnification of the cooling experiments of Figure 2.6 in the temperature range preceding the crystallization of the polymer. The small exotherms are associated to the crystallization of DMDBS. As expected, latent heat of crystallization and peak temperature increase with DMDBS concentration. For clarity, curves are shifted to the same baseline. Table 2.1: Summary of experimental data obtained from DSC data shown in Figure 2.6. DSC DSC Tpeak and Tonset represent peak and onset temperature of the exotherm associated to crystallization of the polymer. tc is the crystallization time defined as DSC DSC DSC tc = (Tonset −Tcompl ) / dT dt where Tcompl corresponds to the completion of the DSC crystallization and dT dt is the cooling rate (=10 ºC/min). Tps represents the peak temperature of the exotherm associated to DMDBS crystallization. X DSC is the degree of crystallinity of the polymer. DSC Tpeak DSC Tonset [ºC] HD120MO DSC Tps X DSC [ºC] ∆H [J·g-1] tc [s] 113 120 95.3 123.5 0.3% DMDBS – B03 131 135 107.7 68.5 149 52 0.7% DMDBS – B07 132 135 107.7 53.6 175 52 1% DMDBS – B1 132 135 103.5 47.3 189 50 [ºC] [%] 46 Upon addition of 0.3 wt% of DMDBS, the crystallization temperature (peak value) of iPP, Tc, increases to 131 ºC. Further addition of DMDBS has nearly no effect on Tc that is 132 ºC for both B07 and B1. Nevertheless, the crystallization peak of the polymer narrows at higher DMDBS contents indicating faster crystallization. Saturation of Tc of iPP with DMDBS concentration was observed also by Kristiansen et al.10, in their data, Tc reaches ~130 ºC at 0.4 wt% DMDBS. Increasing DMDBS concentration, the phase separation occurs 23 Chapter 2 at increasingly higher temperatures. In accordance with WAXD, the final crystallinity of iPP is hardly affected by DMDBS. However, the values measured by DSC, namely 50 %, are noticeably lower than those found with WAXD. Information on the morphology of the system as a function of the temperature is obtained by means of SAXS. Figure 2.8 shows the integrated scattered intensity as a function of the temperature for the neat iPP and the blends with DMDBS. These data can be interpreted in terms of density fluctuations. As expected, in the neat iPP there is no density fluctuation until the polymer starts nucleating at ~120 ºC. While, samples containing DMDBS show more complicated temperature dependence. In fact, when phase separation occurs, DMDBS molecules form crystals denser than the polymer. Figure 2.8: Temperature dependence of the SAXS intensity as a function of DMDBS concentration during cooling at 10 ºC/min and after annealing at 250 ºC for 3 min. In samples containing DMDBS the scattered intensity increases with phase separation because of density fluctuations between DMDBS crystals and the polymer. At lower temperatures, when the polymer crystallizes once again the scattered intensity increases As a consequence, electron density fluctuations are established and the scattered intensity rises to a plateau. At lower temperature, around 135 ºC, independently from DMDBS concentration, nucleation of the polymer triggers a large and abrupt upturn in the intensity. Similar to DSC, some characteristic temperatures for the crystallization of the polymer and of the additive are located and reported in Table 2.2. These data are used to build the phase diagram shown in Figure 2.9 that is used as reference in the rest of this work. 24 Chapter 2 In accordance with Kristiansen et al.10, three different regions, corresponding to three different physical states of the system, are identified: Region I: at high temperatures DMDBS and iPP form a homogenous solution; Region II: at intermediate temperatures the system is phase separated with DMDBS crystallized and iPP molten; Region III: at low temperatures both DMDBS and iPP are crystallized. SAXS Table 2.2: Summary of the SAXS data obtained from Figure 2.9. TcSAXS and Tpeak are, respectively, the onset temperature for polymer crystallization and the SAXS temperature corresponding to the maximum scattered intensity. Tonset ps is the onset SAXS temperature for DMDBS phase separation, and Tplateau is the temperature at which the intensity reaches a constant value (above Tc). SAXS Tplateau TcSAXS SAXS Tpeak SAXS Tonset ps [ºC] [ºC] [ºC] HD120MO 120 [ºC] 108 0.3% DMDBS – B03 135 125 165 150 0.7% DMDBS – B03 135 125 190 175 1% DMDBS – B1 135 127 195 185 Figure 2.9: Phase diagram of the system iPP-DMDBS (from 0 to 1 wt% DMDBS) obtained, on cooling, using SAXS data. Three regions corresponding to three different states can be identified: Region I) homogeneous liquid, Region II) phase separated system with crystallized DMDBS and molten polymer, Region III) both iPP and DMDBS are crystallized 25 Chapter 2 When the polymer crystallizes, in Region III, SAXS allows for the measure of the long period. Figure 2.10 shows the data concerning the neat polymer, B03, B07 and B1 as a function of temperature. As already discussed, the presence of DMDBS leads to an increase in Lp. Figure 2.10: Long period as a function of temperature and DMDBS concentration during temperature ramps with cooling rate of 10 ºC/min. Presence of DMDBS leads to an increase of the long periods that below 80 is quantified in ~4 nm 2.3.3 Morphology of the system in Region II  Two dimensional SAXS images reveal that the increase of the integrated intensity in Region II is caused by an increase of the scattering in all azimuthal directions at low q. Sample images are shown in Figure 2.11. Figure 2.11: SAXS images of the blend B1. Left: material in Region I of the phase diagram and Right: material in Region II of the phase diagram. DMDBS phase separation causes an increase of the scattered intensity in all directions at low q values. For a clear visualization, the scattering of the system in Region I was subtracted. 26 Chapter 2 Such a scattering pattern can be ascribed to the formation of a suspension of randomly oriented DMDBS fibrillar crystals with a length L and a radius R. In this case, the intensity scattered in the q region 2π / L < q < 1 / Rc , can be described with25-27: ⎛ R 2q 2 ⎞ C ⎟ ⎜ ⎟ ⎜ I (q ) = Exp ⎜− c ⎟ ⎟ ⎜ ⎟ q 2 ⎠ ⎝ (2.3) where C is a constant including details on the scatterers like concentration and electron density, while Rc is the radius of gyration of the cross section of the scatterers ( Rc = R / 2 ). From the existing literature, it is known that DMDBS fibrils are basically endless (L ∞), therefore Equation (2.3) is valid for 1/ in this case. Within this limit, Log[I (q ) ⋅ q ] versus q 2 is a straight line with a slope −Rc2 / 2 . Fitting Equation (2.3) to the data points allows for the calculation of Rc and therefore of R. Figure 2.12 provides an example of such a fit demonstrating that a good agreement between experimental data and Equation (2.3) exists for 0.15100s -1 15s / 4s 0.68 2500 0.86 67.8 -1 30s / 2s 1.37 5000 2.1 42.8 60s-1 / 1s 2.73 10000 3.7 41 Similar to what is observed at 142 [%] λ [s] , with the formation of needle-like FIPs, rises quickly to a maximum and then decreases (see Figure 5.7). However, in this case the drop of I SAXS is relatively small and does not correlate with τHMW . Due to the lower D temperature, the early FIPs are more stable than at 142 and their lifetimes t IP τHMW . D Moreover, the drop of I SAXS could also be (partially) explained with the presence of the kebabs that decrease the density fluctuation between the shishes and the melt. 94 Chapter 5 Figure 5.7: Integrated meridional and equatorial SAXS intensities at 137 for 4 s. 15 after a shear of 5.4.5 Separating shish creation from the kebab crystallization  We are now going to show how shishes nucleate the bulk of the polymer. In Chapter 4, the formation of crystalline shishes with shear flow at 142 is discussed. Just like nucleating agents, shishes generated at this high temperature can be used as a heterogeneous substrate for the nucleation of the rest of the molecules at a lower temperature. An exact lattice matching and a good state of dispersion make shishes the perfect nucleating agent. In this paragraph, we discuss the increase in the crystallization temperature and the changes in .  Different flow conditions the morphology induced by shishes generated with shear at 142 are investigated and the cooling starts immediately after flow. Figure 5.8 shows SAXS images taken at 142 shortly after the application of shear. The corresponding De number for HMW molecules are given in Table 5.2. The scattering fingerprint of needle-like FIPs and shishes is clear, independently of shear rate, when γ=100. In all the other cases, in the accessible q range, no changes are observed in the scattering patterns immediately after flow. Table 5.2: Deborah numbers for HMW chains in shear at 142 -1 -1 -1 -1 . 2s 5s 10 s 25 s 50 s-1 100 s-1 HMW Des 0.07 0.2 0.4 0.9 1.8 3.5 HMW Deo 260 650 1300 3250 6500 13000 95 Chapter 5 The scattering patterns of Figure 5.9 suggest the presence of a threshold strain γ for shish formation. Below γ , regardless of , the stretch ratio of the HMW molecules would be insufficient to form shishes. This idea is consistent with the findings of van Meerveld et al.37. However, the absence of an equatorial streak in SAXS does not, necessarily, imply the absence of FIPs and shishes. The equatorial streak is absent also when the concentration of FIPs and shishes is (too) low. This is proven, during cooling, by the growth of lamellae (kebabs) with a high degree of orientation (see Figure 5.9), suggesting the presence of an anisotropic nucleating substrate (FIPs or shishes). According to the simulations of Frenkel and coworkers19, even a single stretched molecule can nucleate parallel lamellae. Surprisingly, in the range of conditions explored in this Chapter, the overall degree of lamellar orientation is found to depend mainly on γ. Very similar orientations are observed at a given γ, independently of . While, much more pronounced differences are observed varying γ (or varying t ) at a given . Figure 5.9 demonstrates that the influence of shear flow on PE morphology is substantial and can be dramatic in some cases. For instance, at γ=100, starting from 25  , isotropic scattering is hardly detectable. For these conditions, all lamellae grow with the c-axis parallel to the shishes (i.e. parallel to the flow direction) yielding a very high level of anisotropy. The emergence of a second order reflection at the meridian of SAXS (see patterns in Figure 5.9) indicates that kebabs are distributed along the flow direction with high regularity. At room temperature, their spacing is 23 nm. Morphology evolution during bulk crystallization is discussed for the case of cooling after a pulse of 100 s-1 for 1 s at 142 , see Figure 5.10. At high temperature, the SAXS equatorial intensity is higher than the meridian and diagonal due to the scattering of FIPs and shishes. During cooling, L dissolve at 142 and D become smaller and, small FIPs, bound to , can crystallize (thermal nucleation) increasing the equatorial intensity. Simultaneously, also the meridian intensity increases but, no maximum can be observed in the intensity profile above T 130 , suggesting uncorrelated objects in the melt. In other words, because of thermal nucleation, small (point-like) FIPs are transformed in stable nuclei. These small FIPs can tumble during shear and, therefore, have random orientations. Moreover, the orientation of these small FIPs is lost already at a relatively small lateral growth. 96 Chapter 5 Figure 5.8: SAXS images few seconds after the application of shear at 142 direction is horizontal. The shear time is indicated in each image. 97 . The flow Chapter 5 Figure 5.9: SAXS images at room temperature after the application of shear at 142 flow direction is horizontal. The shear time is indicated in each image. 98 . The Chapter 5 Figure 5.10: SAXS equatorial and meridional integrated intensities as a function of temperature after shear at 142 (cooling rate is 5 /min). At T 133 , large scale (secondary) nucleation starts. First, the heterogeneous nucleation of kebabs on the surface of shishes causes an increase in I SAXS and, next, at lower temperatures, some isotropic nucleation in the bulk of the sample causes an increase in I SAXS as well. Particularly interesting is the case of mixed morphologies where spherulites coexist with shish-kebabs. In these cases, it is possible to distinguish between the high temperature nucleation of kebabs from the low temperature nucleation of spherulites. An example is the experiment shown in Figure 5.11 (cooling after 10 s for 5 s at 142 ). Figure 5.11: Salient SAXS images during crystallization on cooling after shear at 142 99 . Chapter 5 5.4.6 Crystallization onset temperature after short term shear  As for nucleating agents, the nucleation ability of shishes can be estimated comparing the temperatures where heterogeneous nucleation occurs during cooling. We found that the onset temperature of the kebabs T depends on the flow conditions at 142 conditions described in the previous paragraph, T varies from ~125 . For the flow to ~132 , as reported in Figure 5.12.  Figure 5.12: Onset temperature for the nucleation of kebabs on cooling (5 shear at 142 . Similar to the morphological orientation, T and on t . Consistent with the SAXS data, T /min) after depends mostly on γ rather than on exhibits the highest values at γ=100, where a large amount of shishes (~0.4 %) is formed, and the lowest values at γ=25, where less shishes are formed. The nucleating ability of a substrate can be seen as the product between the quality (efficiency) of the nucleation sites and their number. The quality of nucleation sites is given by matching of the lattice parameters of the substrate and the crystallizing polymer (epitaxy)38-41 and, in this case, is not an issue because of the common unit cell of shishes and kebabs. In addition, it seems reasonable to assume that epitaxy matching does not depend on the flow conditions. In contrast, the quantity of the nucleation sites is determined by aspect ratio ( 1/D ) and concentration of the shishes and both depend on the flow conditions. The individual effects are difficult to separate because a better nucleation ability (higher values of T ) is associated with more nucleation sites that can be achieved both at 100 Chapter 5 the same concentration increasing the aspect ratio and with the same aspect ratio increasing the concentration. 5.4.7 Melting of shish kebabs  Heterogeneous crystallization of kebabs on shishes yields a well organized and oriented morphology. Because of the high nucleation temperature, kebabs grow thicker than ordinary lamellae. For these reasons shish-kebabs exhibit an enhanced thermal stability. To illustrate the melting of shish kebabs, Figure 5.13 shows SAXS intensities as a function of the temperature during heating of a sample crystallized after a shear of 100 s for 1 s at 142 . Figure 5.13: SAXS equatorial and meridional integrated intensities as a function of temperature heating a sample with fully oriented morphology crystallized on cooling after a shear of 100 /1s at 142 . At ~115 , melting of the kebabs (folded chain crystals) starts and it is completed abruptly when they are superheated to ~142 . Shishes, with their extended chain structure, are more stable than spherulites and melt gradually, surviving up to ~157 . Interestingly, immediately after melting of the kebabs, the scattering fingerprint of shishes becomes again visible in the SAXS. This indicates that, if the temperature is not too high, shishes generated with shear at 142 retain their morphology after melting the kebabs. This behavior is 101 Chapter 5 similar to that of some nucleating agents. Kebabs that can melt leaving intact the shishes were already observed and termed macrokebabs by Keller3. During cooling, after heating up to 148 , high temperature nucleation of kebabs is still observed, see Figure 5.14. Figure 5.14: SAXS images during heating and cooling runs around the melting point of the material. This suggests that melting of the shishes begins with losing order at small length scales. As a consequence, chain stretch rapidly vanishes and disrupts the structure of the shishes. However, as suggested by some rheological observations42, the long range order, on the length scale of the whole molecule, can still be retained. Apparently, the long lasting orientation of HMW chains is sufficient for the epitaxial nucleation of kebabs. On heating from room temperature, independently of the shear history at 142 ~142 , kebabs always melt at , i.e. a few degrees higher than spherulites do. In this respect, it is interesting the follow the melting of a sample with a mixed morphology. For instance, Figure 5.15 reports SAXS images of the melting of a sample crystallized after a shear of 5 s 102 for 10 s at 142 . Chapter 5 Figure 5.15: Sequence of SAXS images describing morphological changes during melting of a sample with a fully oriented morphology. At low temperatures, this sample shows a morphology that is a combination of oriented and isotropic distributed lamellae suggesting the coexistence of shish kebabs and spherulites. Upon heating, the isotropic scattering coming from spherulites disappears first at 140 . The meridional lobes coming from kebabs disappear later, at 142 . Shishes were not observed in SAXS during the crystallization of this sample and correspondingly are not observed on heating after melting of the kebabs. 5.5 Conclusions  In this Chapter, we studied the mechanism of formation and melting of shish kebabs in a sheared bimodal HDPE melt. The results indicate that flow induced precursors (FIPs) play a crucial role in structural and morphological developments during crystallization. FIPs are metastable structures with very low or no crystallinity that arise with stretching of the longest chains in the melt. FIPs can be formed at temperatures surprisingly high compared with T and, depending on the conditions, can grow into crystals or dissolve in the melt. When FIPs are formed at very high temperatures, for instance 160 or 200 ,  they are undetectable with SAXS and WAXD but still show strong influence on the crystallization of the polymer. On cooling, FIPs increase the nucleation temperature and promote orientation at structural and morphological levels. The formation of needle-like FIPs and shishes demands stretch of the HMW chains. But this is not the only requirement. Shearing at 142 103 , we Chapter 5 found that the total strain plays a clear role. It seems that a critical strain should be exceeded to obtain a large quantity of shishes. On cooling, needle-like FIPs and shishes formed at 142 act as nucleating agents. They can raise the nucleation temperature to up 132 and template the morphology of the polymer. With time resolved SAXS, it is possible to differentiate between the high temperature epitaxial nucleation of kebabs and the nucleation of lamellae randomly assembled in spherulites, at lower temperatures. We observe that ~0.4 % of shishes is sufficient to produce a fully oriented morphology on cooling. Such a conclusion is consistent with our findings on iPP melts containing fibrils formed from the nucleating agent DMDBS where a similar small amount (~0.5 wt%) of pre-aligned DMDBS fibrils is sufficient to induce, on cooling, a fully oriented polymer morphology. Shish kebabs obtained on cooling after shear at 142 are well organized and this gives them an enhanced thermal stability. At morphological level, melting is a sequence of events symmetrical to crystallization. In fully oriented morphologies, kebabs melt before shishes and yield again SAXS patterns with equatorial scattering only. In mixed morphologies, spherulites melt before kebabs and, similar to the crystallization process, no equatorial scattering is observed. In this Chapter, the transformation of FIPs into extended chain crystals and the nucleating ability of these crystals is addressed. FIPs are partially disordered bundles of molecules with undetectable crystallinity. In fast short term shear, a question arises whether FIPs are generated during or after cessation of flow. In the following Chapter, an attempt has been made to answer this question. 5.6 References  1. 2. 3. 4. 5. Schrauwen, B. A. G.; Breemen, L. C. A. v.; Spoelstra, A. B.; Govaert, L. E.; Peters, G. W. M.; Meijer, H. E. H. Macromolecules 2004, 37, 8618-8633. Gahleitner, M.; Wolfschwenger, J.; Bachner, C.; Bernreitner, K.; Neiβl, W. Journal of Applied Polymer Science 1996, 61, 649-657. Keller, A.; Kolnaar, H. W. H., Flow induced orientation and structure formation. VCH: New York, 1997; Vol. 18. Meijer, H. E. H.; Govaert, L. E. Prog. Polym. Sci. 2005, 30, 915-938. Ward, I. M.; Sweeney, J., An introduction to the mechanical properties of solid polymers. John Wiley & Sons: Weinheim, 2004. 104 Chapter 5 6. 7. 8. 9. 10. 11. 12. 13. 14. 15. 16. 17. 18. 19. 20. 21. 22. 23. 24. 25. 26. 27. 28. 29. 30. 31. 32. 33. 34. 35. Kumaraswamy, G.; Issaian, A. M.; Kornfield, J. A. Macromolecules 1999, 32, 75377547. Janeschitz-Kriegl, H.; Ratajski, E.; Stadlbauer, M. Rheol Acta 2003, 42, 355. Devaux, N.; Monasse, B.; Haudin, J. M.; Moldenaers, P.; Vermant, J. Rheol. Acta 2004, 43, 210-222. Bashir, Z.; Odell, J. A.; Keller, A. Journal of Materials Science 1984, 19, 3713-3725. Hill, M. J.; Barham, P. J.; Keller, A. Colloid & Polymer Sci. 1980, 258, 1023-1037. Barham, P. J.; Keller, A. Journal of Materials Science 1985, 20, 2281-2302. Hsiao, B. S.; Yang, L.; Somani, R. H.; Avila-Orta, C. A.; Zhu, L. Physical Review Letters 2005, 94, 117802. Smith, P.; Lemstra, P. J.; Pijpers, J. P. L.; Kiel , A. M. Colloid & Polymer Sci. 1981, 259, 1070-1080. Hsieh, Y.; Ju, J. Journal of Applied Polymer Science 1994, 53, 347-354. Loos, J.; Schimanski, T.; Hofman, J.; Peijs, T.; Lemstra, P. J. Polymer 2001, 42, 3827-3834. Liu, T.; Lieberwirth, I.; Petermann, J. Macromol. Chem. Phys. 2001, 202, (14), 29212925. Pennings, J.; Kiel , A. M. Kolloid ZZ Polym 1965, 205, 160. Binsbergen, F. L. Nature 1966, 211, 516-517. Hu, W.; Frenkel, D.; Mathot, V. B. F. Macromolecules 2002, 35, 7172-7174. Dukovski, I.; Muthukumar, M. Journal of Chemical Physics 2003, 118, (14), 66486655. Keum, J. K.; Burger, C.; Zuo, F.; Hsiao, B. S. Polymer 2007, 48, 4511-4519. Zuo, F.; Keum, J. K.; Yang, L.; Somani, R. H.; Hsiao, B. S. Macromolecules 2006, 39, 2209-2218. Kumaraswamy, G.; Kornfield, J. A.; Yeh, F.; Hsiao, B. S. Macromolecules 2002, 35, 1762-1769. Ogino, Y.; Fukushima, H.; Matsuba, G.; Takahashi, N.; Nishida, K.; Kanaya, T. Polymer 2006, 47, 5669-5677. Keum, J. K.; Zuo, F.; Hsiao, B. S. Journal of Applied Crystallography 2007, 40, 4851. Seki, M.; Thurman, D. W.; Oberhauser, J. P.; Kornfield, J. A. Macromolecules 2002, 35, 2583-2594. Nogales, A.; Hsiao, B. S.; Somani, R. H.; Srinivas, S.; Tsou, A. H.; Balta-Calleja, F. J.; Ezquerra, T. Polymer 2001, 42, 5247-5256. Yang, L.; Somani, R. H.; Sics, I.; Hsiao, B. S.; Kolb, R.; Lohse, D. J. Phys. Condens. Matter 2006, 18, 2421-2436. Yang, L.; Somani, R. H.; Sics, I.; Hsiao, B. S.; Kolb, R.; Fruitwala, H.; Ong, C. Macromolecules 2004, 37, 4845-4859. Macosko, C., Rheology : principles, measurements, and applications. VCH: Weinheim, 1994. Balta-Calleja, F. J.; Vonk, C. G., X-ray scattering of synthetic polymers. Elsevier: Amsterdam, 1989. Wilchinsky, Z. W., Advances in X-ray analysis. 1962; Vol. 6. Wunderlich, B., Thermal analysis of polymeric materials. Springer-Verlag: Berlin, 2005. Kimata, S.; Sakurai, T.; Nozue, Y.; Kasahara, T.; Yamaguchi, N.; Karino, T.; Shibayama, M.; Kornfield, J. A. Science 2007, 316, (5827), 1014 - 1017. Coppola, S.; Balzano, L.; Gioffredi, E.; Maffettone, P. L.; Grizzuti, N. Polymer 2004, 45, 3249-3256. 105 Chapter 5 36. 37. 38. 39. 40. 41. 42. Somani, R. H.; Yang, L.; Hsiao, B. S. Physica A 2002, 304, 145-157. van Meerveld, J.; Peters, G. W. M.; Hutters, M. Rheol Acta 2004, 44, 119-134. Wittmann, J. C.; Lotz, B. J. Polym. Sci., Polym. Phys. Ed. 1981, 19, 1837. Thierry, A.; Straupe, C.; J., W.; Lotz, B. Macromol Symp 2006, 241, 103-110. Fillon, B.; Thierry, A.; Lotz, B.; Wittmann, J. C. Journal of Thermal Analysis 1994, 42, 721-731. Lotz, B.; Wittmann, J. C. J. Polym. Sci., Polym. Phys. Ed. 1986, 24, 1559. Bent, J.; Hutchings, L. R.; Richards, R. W.; Gough, T.; Spares, R.; Coates, P. D.; Grillo, I.; Harlen, O. G.; Read, D. J.; Graham, R. S.; Likhtman, A. E.; Groves, D. J.; Nicholson, T. M.; McLeish, T. C. B. Science 2003, 19, 1691 - 1695. 106 Chapter 6*  Metastable structures during fast short term  shear  The goal of this Chapter is to study the crystallization of isotactic polypropylene (iPP) during and immediately after a strong pulse of shear (i.e. with a high Deborah number) on the undercooled melt. This is possible using high flux synchrotron X-ray scattering (SAXS and WAXD) that allows for a high frame rate (3 frames/s) maintaining a good signal to noise ratio. We found that shear rate is the dominating parameter for structure formation during shear. When the shear rate is high enough, crystals with a high degree of orientation in the shear direction can be formed already during shear. In contrast, for lower shear rates, no crystalline structures were observed during shear. However, SAXS shows an equatorial streak indicating needle-like scatterers aligned with the flow direction. These scatterers are metastable precursors of crystallization. Although these precursors are not crystalline in the early stages, they crystallize after cessation of the flow and template the nucleation of lamellae with caxis parallel to the flow direction (shish-kebab morphology). 6.1 Introduction  Bulk crystallinity1-5 affects many of the physical properties of semicrystalline polymeric materials6-12. Universal rules are difficult to establish as these materials exhibit a processing-property relation, i.e. the final properties of the material are affected by the processing conditions. Temperature history and, especially, flow alter the crystallization behavior. Flow can increase the crystallization rate by decades compared to quiescent conditions13-18 and, in addition, the crystalline morphology can change dramatically19-24. In quiescent or quasi-quiescent conditions morphology is dominated by spherulites25, three dimensional assemblies of randomly oriented folded chain lamellae. For strong enough flows, spherulites are replaced by shish-kebabs, composite crystallites with an extended chain fibrillar core (shish) dressed with disk-like folded chain lamellae (kebabs)26-32. The origin of * Partially reproduced from: L. Balzano et al. ‘Metastable structures during fast short term shear’, Macromolecules 2008 (Submitted) 107 Chapter 6 this morphology is the topic of a long standing discussion31, 33-37 that begun in the 1960s. In Chapters 4 and 5, some of the classic ideas are combined with new observations, to propose a mechanism for the formation of shish-kebabs that starts with stretch of the longest chains and passes through metastable precursors with undetectable crystallinity. In short term shear, after cessation of the flow, a selection takes place among the precursors. Those exceeding some critical dimension transform into crystalline structures, the remainder dissolve back into the melt state. The ratio between crystallizing and dissolving precursors and the rate of crystallization and dissolution are determined by the temperature. A real competition between crystallization and dissolution of flow induced precursors (FIPs) can be observed only at high enough temperatures. For instance, in the case of polyethylene (PE), it is necessary to be in 0 proximity to or even above the equilibrium melting point Tm0 ( Tm PE = 141.2 °C ). Although there is no undercooling Δ T , this does not contradict thermodynamics. In fact, flow raises the effective melting temperature by reducing the entropy and, thus, generates the undercooling ΔT f that is indispensible for nucleation13, 38. At lower temperatures, the driving force due to flow contribute to make most of the precursors stable from the moment of their creation33. Here, dissolution is absent or limited to a small fraction of precursors. In many of the experiments available in literature, the study of FIPs is based on short term flow protocols33, 39-44 where, a pulse of flow is applied for few seconds, and the evolution of the system is studied after cessation of the flow. In this scenario, some questions arise: when are actually FIPs created and when do they crystallize? During or after cessation of the flow? Reminding that FIPs are disordered bundles of molecules held together only by weak interaction forces, another question that arises is: are these interaction forces strong enough to prevent flow from destroying FIPs and allow them to crystallize during flow? Giving answers to these questions is the aim of this Chapter. In a short term flow protocol, flow lasts, typically, only for few seconds and this restricts the sampling time to a fraction of a second. For this reason, our investigation demands the use of synchrotron X-ray scattering with a high flux of photons. The material investigated is isotactic polypropylene (iPP). 108 Chapter 6 6.2 Materials and methods  6.2.1 Materials  The polymer used in this work is a commercial homopolymer grade of isotactic polypropylene (iPP). The material, labeled 15M10, was obtained by DSM (Geleen, The Netherlands) and contained no other additives than stabilizers. The specifications of the material are reported in Table 6.1. Table 6.1: Specification of the iPP used in this work. Material Mw [kg/mol] Mw/Mn [%mmmm] Tm [°C] 15M10 350 5.6 96.2 161 6.2.2 X‐ray characterization  Small angle X-ray scattering (SAXS) was performed at the beamline ID02 of the European Synchrotron Radiation Facility (ESRF,Grenoble). The samples were irradiated with a wavelength λ=0.995 Å and two dimensional images were recorded using a Frelon detector, with a resolution of 1024x1024 pixels and a pixel size of 164 µm, placed at 6.5 m from the sample. After subtraction of the scattering of the empty sample holder, images were integrated to obtain the scattered intensity (I) as a function of the modulus of the scattering vector q = ( 4π / λ ) sin(θ / 2) where 2θ is the scattering angle45, 46. The total intensity I SAXS as a function of time was obtained integrating I(q) over the whole accessible q range: I SAXS (t ) = qmax ∫ I (q; t )dq . qmin Wide angle X-ray scattering (WAXS or WAXD) was performed at the beamline ID11 of the ESRF. Experiments were carried out with a wavelength λ=0.508 Å and a sample to detector distance of about 32cm. The images were recorded with a two dimensional Frelon detector with a resolution of 512x512 pixels (pixel size 190  µm). After correction for spatial 109 Chapter 6 distortions and scattering of the empty sample holder, the images were integrated to obtain the intensity distribution I as a function of the scattering angle 2θ . These one-dimensional profiles were used to calculate crystallinity. For this purpose, the scattering of the amorphous component (IA) underneath crystalline peaks (Ic) was approximated with a straight line (see Figure 6.1). Figure 6.1: Example of the procedure to separate crystalline and amorphous scattering. The Miller indices of the scattering reflections are indicated as well. This simplified procedure is expected to give trustworthy results at low crystallinities (up to ~15%) 47. Crystallinity (x) was determined as: x = 100 ⋅ IC IC + I A (6.1) WAXD data were also used for calculating the Hermans’ orientation factors FH of the crystalline c-axis: FH = 3 cos 2 β − 1 2 110 (6.2) Chapter 6 where β is the azimuthal angle. For iPP, where a 00l reflection is not observed, cos 2 β can be obtained with Wilchinsky’s formula48. Using geometrical relations, the orientation of the c-axis is calculated from the orientation of the 110 and 040 crystal planes: cos 2 β = 1 − 0.901 cos 2 β 040 − 1.099 cos 2 β110 (6.3) with the definition: π /2 cos β = 2 ∫ I (β ) ⋅ cos 0 2 β ⋅ sin β ⋅ d β π /2 (6.4) ∫ I (β ) ⋅ sin β ⋅ d β 0 6.2.3 Shear experiments  Shear experiments were performed in combination with SAXS and WAXD using a Linkam CSS-450 Shear Cell where, to avoid unwanted scattering, the original glass plates were replaced with kapton. The setup is shown in Figure 6.2. Figure 6.2: Schematic drawing of the shear device. When shearing, a metal spoke of the rotating plate lies periodically in front of the incoming beam, the corresponding (dark) scattering images are discarded from analysis. 111 Chapter 6 6.3 Results and Discussion  6.3.1 Flow conditions in short term shear  It is well established that shear flow promotes orientation and stretch of polymer molecules in the melt23, 24, 49-51 . If τ D and τS are the relaxation times for orientation (disengagement time) and stretch (chain retraction time) of the longest molecules, orientation of these molecules is achieved when the Deborah number for orientation DeO = τ D γ exceeds unity and stretch when the Deborah number for stretch DeS = τS γ exceeds unity49. The ratio τ D / τ S is similar to the number of entanglements per chain Z 52-54 . For the iPP under consideration, Z 100 and, therefore, DeO ≅ 100DeS . In other words, chain orientation is attained at shear rates nearly two decades lower than chain stretch. In this Chapter, we aim to look at crystallization under shear conditions providing orientation and stretch to the molecules, therefore, our experimental conditions are chosen such that DeS >>1. A high DeO is consequently achieved as well. A useful way to study the flow induced crystallization (FIC) of a polymer melt is short term shear. In this experiment, after annealing at high temperature, the melt is brought to the desired test temperature where shear is applied for a limited time, typically few seconds, and the subsequent crystallization of the polymer is observed in absence of flow. By choosing a short shear time it is hoped that flow effects are separated from crystallization effects, i.e. no noticeable material changes occur during the flow and thus the ‘normal’ rheological behavior can be used to characterize the flow strength ( DeO and DeS ). The flow conditions are determined by three parameters: temperature, shear rate ( γ ) and shear time ( ts ). Once these are assigned, the total strain ( γ = γ ⋅ ts ) is also fixed. A fair comparison between different experiments is possible only when holding γ constant. 112 Chapter 6 6.3.2 Flow induced precursors during short term shear  To generate FIPs, we select shear conditions with a relatively high strain value, γ = 180 . This allows for spanning a broad range flow conditions, from relatively slow and long to relatively fast and short. FIPs originate from stretching of the longest chains in the melt 23, 28, 41, 44, 55 , therefore we restrict γ such as DeS >> 1 . For the iPP considered, τ D ≅ 100 s and τ S ≅ 1 s at 145 °C 56 . Therefore, shear rates of 60 s-1, 90 s-1 and 180 s-1 suite our experimental requirements. The shear time is varied accordingly to keep the total strain γ = 180 . Table 6.2 gives an overview of the selected shear conditions. WAXD patterns acquired at a rate of 3 frames/s during and immediately after short term shear at 145 °C are shown in Figure 6.4. In these figures, as in the rest of the Chapter, the time t=0 corresponds to cessation of the flow. During flow (t 1 ). As revealed by the appearance of SAXS meridional lobes, after cessation of the flow, FIPs start crystallizing and serve as a heterogeneous substrate for the nucleation of stack of lamellae with c-axis parallel to the flow direction. Figure 6.3: SAXS patterns during and immediately after short term shear at 145 °C with 60s1 /3s. 6.3.3 Crystallization after short term shear.  After cessation of shear (t=0), samples are allowed to crystallize isothermally at 145 SAXS SAXS °C. Figure 6.5 reports the SAXS intensities in the equatorial ( I Eq ) and meridional ( I Mer ) 114 Chapter 6 Figure 6.4: Equatorial region of WAXD patterns during and after short term shear at 145°C with γ =180. Flow direction is vertical. The arrows indicate arched crystalline reflections. 115 Chapter 6 SAXS SAXS regions for the experiment 60s-1/3s. At the beginning, I Eq is much higher than I Mer SAXS because of the equatorial streak coming from needle like FIPs. However, while I Eq SAXS remains nearly constant, I Mer grows quickly and after few seconds the situation is reversed, SAXS SAXS SAXS i.e. I Mer > I Eq . The growth of I Mer is a consequence of the onset of meridional lobes associated with the decoration of shishes and FIPs with stacks of lamellae (kebabs) with the c-axis oriented with the flow direction. The formation of a similar shish-kebab morphology is observed also for 90s-1/2s and 180s-1/1s after cessation of flow. Figure 6.5: SAXS meridional and equatorial intensities at 145 °C as a function of time after the application of shear (60s-1/3s). As shown in Figure 6.6, the growth of kebabs raises the crystallinity and, simultaneously, the crystalline orientation factors (Figure 6.7) tend to decrease. This can be explained with the homogeneous nucleation of randomly oriented lamellae in the bulk of the sample and/or with bending of the kebabs after a certain diameter is exceeded (see Figure 6.8)61. In the experiments 90s-1/2s and 180s-1/1s, where crystallization is observed already during shear, the orientation is significantly higher than in 60s-1/3s. As expected, the data of Figure 6.6 indicate that an increase in γ is more efficient to enhance both nucleation and orientation than an increase in ts . 116 Chapter 6 Figure 6.6: Crystallinity developing at 145 °C after the application of 180 strain units. Figure 6.7: Orientation developing at 145 °C after the application of 180 strain units. 117 Chapter 6 Figure 6.8: Bending of kebabs occurring after a certain diameter is exceeded. More details on structure formation immediately after cessation of shear can be obtained defining a degree of space filling after flow: Φ(t ) = x(t ) − x0 x∞ − x0 (6.5) x0 is the crystallinity already present at the time t=0 and, therefore, generated during flow, while, x∞ is the saturation level for the crystallinity, attained at long times. Note that the definition (6.5) is based on the assumption that, at long times, a complete space filling ( Φ =1) is attained despite the relatively low degree of crystallinity ( x∞ ). Φ(t ) can be described with the Avrami equation62-64: Φ(t ) = 1 − exp(−kt n ) (6.6) where k is a rate constant and n the Avrami exponent that indicates the geometry of the growing crystallites. Equation (6.6) can be re-written as: Ln {− Ln [1 − Φ (t ) ]} = Ln ( k ) + n ⋅ Ln (t ) (6.7) In the early stages, Ln {− Ln [1 − Φ (t ) ]} as a function of Ln (t ) (the so called Avrami plot) is a straight line with slope n. The Avrami plot, for the crystallization following the cessation of the shear, is shown in Figure 6.9. Fitting Equation (6.7) on the data yields the parameters given in Table 6.3. 118 Chapter 6 Table 6.3: Parameters used for fitting Equation (6.6) onto crystallization data of Figure 6.6. x∞ [%] x0 [%] n k [s-n] 60 s-1/3s 14.5 0 1.34 2.70·10-3 90 s-1/2s 16.2 0.1 1.37 6.78·10-3 180 s-1/1s 18.8 2 1.38 8.20·10-3 Figure 6.9: Avrami plot constructed with crystallinity data after short term shear at 145 °C. Independent of the flow conditions, n 1.36±0.02 indicates that, in all cases, space filling occurs because of the athermal growth of circular lamellae (kebabs). The only difference is the rate constant k that increases with γ , in line with the increased number of kebabs at higher shear rates and consistent with an increased number of shishes as well. Combining equations (6.5) and (6.6), it is possible to obtain a description of the crystallinity as a function of time: x(t ) = x0 + ( x∞ − x0 ) ⎡1 − exp ( −kt n ) ⎤ ⎣ ⎦ (6.8) As expected, using the fitting parameters given in Table 6.3, a good agreement with the measured data is found (lines in Figure 6.6). Some remarkable observations emerge from this analysis: 119 Chapter 6 • After a short term shear, a good description of the crystallization kinetics can be obtained with the Avrami equation. To this end, it is necessary to define the space filling Φ(t ) accounting for the crystallinity generated already during flow ( x0 in Equation (6.5)). Although this value can be very small, it has large influence in the early stages. Furthermore, a complete space filling at long times needs to be hypothesized, despite the relatively low crystallinity. • The data of Figure 6.6 seem to indicate the presence of an ‘induction time’ for the growth of kebabs, even for 180s-1/1s where a crystallinity of 2 % is attained already during flow. However, although in all experiments there is a clear characteristic time at which x starts increasing massively, this induction period should not be considered as a time where structure formation is frozenin. Figure 6.9 indicates that, in all cases, starting form the first measurement (t=0.3 s), major changes take place in the space filling. • At long times, in isothermal conditions, crystallinity saturates to a plateau level ( x∞ ) that depends on the flow conditions. Higher γ yield higher final crystallinity. This behavior can be interpreted in terms of driving force for crystallization, i.e. in terms of Gibbs free energy (G) difference between the melt and the crystal. Flow increases the free energy of the melt generating more driving force for nucleation (see Figure 6.10). Figure 6.10: Effect of flow on the free energy. 120 Chapter 6 6.4 Conclusions  We studied the early stages of the flow induced crystallization of iPP, during and immediately after the application of a strong (high De) shear pulse. At 145 °C, with a total strain γ = 180 , X-ray scattering data (SAXS and WAXD) suggest that shear rate is the dominant parameter for structure development during flow. When the shear rate is high enough (> 90 s-1) crystalline structures can be formed already during shear. In these conditions, WAXD suggests the presence of crystallinities up to 2% and a high degree of orientation of the crystals with the flow direction. In contrast, at lower shear rates (for instance 60 s-1) no sign of crystalline structures is observed during shear. However, simultaneously, SAXS indicates the presence of needle like scatterers aligned with the shear direction. As soon as the flow stops, the crystalline or non-crystalline structures generated during shear act as a nucleating substrate for the heterogeneous formation of kebabs. The crystalline morphology is, thus, strongly tied with the structures present in the early stages. 6.5 References  1. 2. 3. 4. 5. 6. 7. 8. 9. 10. 11. 12. 13. 14. 15. Strobl, G., The physics of polymers. Springer: Berlin, 1997. Wunderlich, B., Thermal analysis of polymeric materials. Springer-Verlag: Berlin, 2005. Keller, A. Phyl. Mag. 1957, 2, 1171. Lotz, B. Journal of Macromolecular Science, Part B-Physics 2002, B41, 685-709. Natta, G. J. Polym. Sci. 1955, 16, 143. Schrauwen, B. A. G.; Breemen, L. C. A. v.; Spoelstra, A. B.; Govaert, L. E.; Peters, G. W. M.; Meijer, H. E. H. Macromolecules 2004, 37, 8618-8633. Kristiansen, M.; Werner, M.; Tervoort, T.; Smith, P.; Blomehofer, M.; Schmidt, H. W. Macromolecules 2003, (36), 5150-5156. Smith, P.; Lemstra, P. J. Journal of Materials Science 1980, 15, 505-514. 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Macromolecules 2001, 34, (14), 5030 5036. Azzurri, F.; Alfonso, G. C. Macromolecules 2005, 38, 1723-1728. Keum, J. K.; Zuo, F.; Hsiao, B. S. Journal of Applied Crystallography 2007, 40, 4851. Jerschow, P.; Janeschitz-Kriegl, H. International Polymer Processing 1997, 12, (1), 72-77. Van der Beek, M. H. E.; Peters, G. W. M.; Meijer, H. E. H. Macromolecules 2006, 39, (5), 1805 -1814. Pogodina, N. V.; Winter, H. H.; Srinivas, S. Journal of Polymer Science: Part B: Polymer Physics 1999, 37, 3512-3519. Elmoumni, A.; Gonzalez-Ruiz, R. A.; Coughlin, E. B.; Winter, H. H. Journal of Chemical Physics 2005, 206, 125-134. Stribeck, N., X-ray scattering of soft matter. Springer Laboratory: 2007. 122 Chapter 6 46. 47. 48. 49. 50. 51. 52. 53. 54. 55. 56. 57. 58. 59. 60. 61. 62. 63. 64. Balta-Calleja, F. J.; Vonk, C. G., X-ray scattering of synthetic polymers. Elsevier: Amsterdam, 1989. Kumaraswamy, G.; Verma, R. K.; Kornfield, J. A.; Yeh, F.; Hsiao, B. S. Macromolecules 2004, 37, (24), 9005 -9017. Wilchinsky, Z. W., Advances in X-ray analysis. 1962; Vol. 6. van Meerveld, J.; Peters, G. W. M.; Hutters, M. Rheol Acta 2004, 44, 119-134. Dukovski, I.; Muthukumar, M. Journal of Chemical Physics 2003, 118, (14), 66486655. Somani, R. H.; Yang, L.; Hsiao, B. S. Physica A 2002, 304, 145-157. Dealy, J. M.; Larson, R. G., Structure and Rheology of Molten Polymers. Hanser Gardner Pubns: Cincinnati, 2006. Doi, M.; Edwards, S. F., The theory of polymer dynamics. Clarendon Press: Oxford, 1986. McLeish, T. C. B. Advances in Physics 2002, 51, (6), 1379-1527. Yang, L.; Somani, R. H.; Sics, I.; Hsiao, B. S.; Kolb, R.; Lohse, D. J. Phys. Condens. Matter 2006, 18, 2421-2436. Peters, G. W. M.; Swartjes, F. H. M.; Meijer, H. E. H. Macromol. Symp. 2002, 185, 277-292. Fujiyama, M.; Wakino, T. Journal of Applied Polymer Science 1988, 35, 29-49. Kolb, R.; Seifert, S.; Stribeck, N.; Zachmann, H. G. Polymer 2000, 41, 1497-1505. Ran, S.; Zong, X.; Fang, D.; Hsiao, B. S.; Chu, B. Macromolecules 2001, 34, 25692578. Ruland, W. Journal of Polymer Science: Part C 1969, 28, 143-151. Xu, J.; Johnson, M.; Wilkes, G. L. Polymer 2004, 45, (15), 5327-5340. Avrami, M. J. Chem. Phys. 1939, 7, 1103-1112. Avrami, M. J. Chem. Phys. 1940, 8, 212-224. Avrami, M. J. Chem. Phys. 1941, 9, 177-184. 123 Chapter 6 124 Chapter 7  Conclusions and recommendations  7.1 Conclusions / Technology assessment  Polyolefins, basically polyethylene (PE) and polypropylenes (PP), are materials with a vast number of applications. However, nearly seventy years after the industrial launch, their potential is not yet fully exploited. Smart processing, additives and control over molecular parameters can still enhance the performances of these olefin-based materials. The mechanical, optical and transport properties of polyolefins are strictly related to crystallinity and crystalline morphology. For this reason, one of the key issues is understanding the crystallization behavior, especially in presence of flow, in order to be able to direct it in the desired way. Flow can enhance the crystallization kinetics, can change the morphology and can make crystallization happen even at temperatures where it is not expected, i.e. above the melting point (extended chain crystals). During processing, the flow conditions (geometry, temperature, flow rate) should be tailored in primis to the molecular weight and to the molecular weight distribution of the melt but also to the additives that are present. Self-assembling sorbitol-based additives (DBS, MDBS, DMDBS, etc.) are often used as nucleating/clarifying agent for isotactic polypropylene (iPP). At concentrations between 0.1 and 1 wt%, these additives can crystallize before iPP, leading to an interconnected network of nano-fibrils. The polymer crystallizes forming disk-like lamellae that grow in the direction of the radius of these nanofibrils. As a consequence, the orientation of iPP lamellae is controlled by the orientation of the nanofibrils at the onset of iPP crystallization. In absence of flow, DMDBS nanofibrils grow randomly oriented in the melt but they can be aligned with flow. Like in many other cases, flow can switch lamellar orientation on and off. However, when sorbitol based additives are used, there is a rather sharp transition between the flow conditions that can yield orientation and those that cannot. This transition is the crystallization of the additive. When a random lamellar orientation is desired, flow should be 125 Chapter 7 applied at temperatures higher than the crystallization of the additive. On the other hand, when a high degree of lamellar orientation is desired, flow should be applied at temperatures lower than the crystallization of the additive. Such a procedure, in combination with ~0.5 wt% of DMDBS, can lead to a fully oriented iPP morphology after the application of flow at a temperature as high as ~200 °C. Heterogeneous nucleation on oriented substrates can be used in several processes (e.g. spinning) where a high degree of lamellar orientation in the final product is demanded. The outstanding performances of sorbitol-based additives in iPP are due to: 1) a good dispersion in the polymer matrix; 2) a high aspect ratio of the nanofibrils; 3) a good epitaxy matching. Similar conditions, with a pre-aligned nano-fibrillar substrate templating the morphology of the bulk of the polymer, can be generated also in polyethylene (PE). Here the nanofibrils are made out of the same polymer and can be produced in-situ by inducing, with flow, the extended chain crystallization of the longest molecules of the melt. The molecular weight distribution of the melt needs to be carefully tailored on the flow conditions. Very broad and bimodal molecular weight distributions can fulfill the requirements as demonstrated in the thesis. 7.2 Recommendations for future research  Like all scientific contributions, this thesis tries to give few answers and puts forward many questions. Several aspects demand further clarification. Two outstanding questions are: • Heterogeneous crystallization is based on the epitaxial relation (lattice matching) between the nucleating agent and the crystallizing polymer. Polymer molecules prefer to crystallize onto heterogeneous substrates because, below the melting point, the contact (adsorption) of a polymer molecule onto the foreign particle minimizes the energy of the system. What happens above the melting point? Is this scenario retained? And what are the consequences on the rheology of the melt? • Formation of needle-like flow induced precursors with undetectable crystallinity is a typical example of ‘SAXS before WAXD’. Flow induced precursors assume the form of structures with only short range order. In this 126 Chapter 7 thesis, by calling them precursors, we confine them in a black box. What is the structure of flow induced precursors? Do they differ from the precursors in quiescent conditions? Do they alter the rheology of the system? Do chains need to disentangle to generate a FIP? • Polymers with broad molar mass distributions are often used in applications (e.g. film blowing). In these melts, there is a subtle balance between the relaxation times of long and short chains based on the life time of entanglements that can be very limited (constraint release). For the formation of oriented morphologies, what is more important? The high or the low molar mass? 127 Chapter 7 128      Samenvatting  Synthetische polymeren worden momenteel geproduceerd in grote hoeveelheden, meer dan 200 miljoen ton op jaarbasis. Meer dan 70% van dit produktie volume wordt ingenomen door de zogenaamde klasse van “Commodity Plastics”, te weten de bekende vier bulkpolymeren: polyethyleen (PE), polypropyleen (PP), polystyreen (PS) en polyvinylchloride (PVC). Polyethyleen en Polypropyleen vormen samen de subgroep Polyolefinen en deze klasse van polymeren is met meer dan 100 miljoen ton op jaarbasis de belangrijkste groep van industriële polymeren. De aanduiding Polyolefinen of PE en PP is generiek, bijv. er worden van PE en PP tal van soorten geproduceerd die onderling verschillen in molmassa (ketenlengte), ketenstructuur (stereoregulariteit, vertakkingen etc.), allemaal toegesneden op specifieke toepassingen. De eigenschappen van synthetische polymeren worden niet alleen bepaald door de chemische structuur van het polymeer maar ook door de verwerkingscondities tot eind produkt. Ketenoriëntatie is met name bepalend voor de mechanische eigenschappen met als extreem voorbeeld het verschil in eigenschappen tussen flexibele folies en containers en de supersterke vezel Dyneema® van DSM, allemaal gebaseerd op (lineair) PE maar met een verschil in keten oriëntatie en molmassa. Naast de chemische structuur en verwerkingscondities kan men de eigenschappen van polymeren ook sterk beïnvloeden door het gebruik van additieven en er zijn tal van additieven ontwikkeld zoals kleurstoffen, brandvertragers, anti-oxydanten en om bijv. de vloei van de polymeren te verbeteren bij verwerking of ter verbetering van de mechanische eigenschappen zoals glasvezels ter verhoging van de stijfheid. In feite is een plastic (compound) een polymeer + additief! In dit proefschrift is aandacht besteed aan het phenomeen van nucleatie, het begin van het kristallisatie-proces. Polymeren zoals PE en PP kunnen kristalliseren en het optreden van kristallisatie is zeer belangrijk voor de mechanische eigenschappen. Zonder kristallisatie zouden polymeren zoals PE en PP rond kamertemperatuur en/of verhoogde temperatuur niet gebruikt kunnen worden omdat het dan slechts hoog-visceuze vloeistoffen zouden zijn. Kristallisatie c.q. ordening van de lange polymeer moleculen in de bekende zogenaamde gevouwen-keten kristallen (lamellar folded-chain crystals) geeft versteviging aan PE en PP 129 tot aan het smeltpunt. Polypropeen (PP) is een uiterst traag kristalliserend polymeer en dat is inherent aan de structuur van het molecuul, een 31 helix in de kristallijne fase. Bij vormgeving, bijv. spuitgieten, is dit kristallisatie proces te traag. Om de overall kristallisatiesnelheid te verhogen worden zogenaamde kiemvormers toegevoegd die het ontstaan van kiemen, resulterend in kristallen, in het materiaal sterk bevorderen. Er zijn twee typen kiemvormers bestudeerd, te weten: - De bekende sorbitol derivaten die commercieel worden gebruikt en die als eigenschap hebben dat ze, gemengd met PP, assembleren in nano-fibrillen bij afkoelen voor dat de PP fase kristalliseert. Op deze nano-fibrillen kan PP gemakkelijk ontkiemen; - Een mengsel van hoog- en laag-moleculair PE, gemaakt via een unieke zogenaamde “one pot” synthese waardoor de beide componenten op moleculaire schaal zijn gemengd. Onder afschuiving, zoals optreedt bij verwerking via de polymere smelt, zal de hoog-moleculaire component gemakkelijker orienteren en deze gestrekte ketens fungeren als kiem voor kristallisatie van de laag-moleculaire matrix PE. Bij het PP/sorbitol systeem werd gevonden dat bij zeer lage concentraties, < 1%, van het gebruikte type 1,3:2,4-bis (3,4-dimethylbenzylideen) sorbitol (DMDBS) zeer efficiënte kiemvorming optreedt en resulteert in zeer kleine PP kristallieten die nagenoeg geen licht verstrooien, dus geeft een transparant PP produkt. DMDBS is oplosbaar in gesmolten PP maar kristalliseert uit in nano-fibrillen, door zelf-assemblage, die fungeren als kiem voor kristallisatie van PP. Door gebruik te maken van gecombineerde temperatuur en afschuifprofielen kunnen de DMDBS nano-fibrillen worden georiënteerd in een bepaalde richting en dit biedt de mogelijkheid om unieke structuren te maken. Bij het PE/PE systeem, respectievelijk laag- en hoog-moleculair PE, werd gevonden dat de hoog-moleculaire component kan worden georiënteerd boven het smeltpunt van de matrix, de laag-moleculaire component. Dit biedt de mogelijkheid tot nucleatie van de laagmoleculaire component op de gestrekte ketens van hoog-moleculair PE., in feite kristalliseert de laag-moleculaire PE component op hetzelfde materiaal, een perfecte vorm van epitaxy! 130 In conclusie, er werden twee systemen bestudeerd die ogenschijnlijk niets met elkaar te maken hebben maar het samenbindend element is het genereren van georiënteerde kiemen resulterend in unieke structuren c.s. morphologie van respectievelijk PP en PE. 131 132 Acknowledgements  Many people have contributed to this thesis and many have contributed to make my last four years in Eindhoven a real good time, I want to express my gratitude to everyone. First of all, I would like to acknowledge the people who offered me the opportunity of doing a PhD and provided the necessary resources: Prof. Piet Lemstra, Prof. Sanjay Rastogi and Dr. Gerrit Peters. I am grateful to Prof. Piet Lemstra for the motivations, the teachings and for instilling in my mind his research style, deeply scientific but always with an eye to applications. This work could not be done without Prof. Sanjay Rastogi, a daily source of inspiration. Thank you, Sanjay, for all the discussions and for the stimulating ideas that led to this thesis. Working with you, I learned a lot, challenging or validating the existing ideas, always with determination and enthusiasm. In addition, I want to thank you for the friendship that made me always feel comfortable. This thesis is largely due to Dr. Gerrit Peters. Hartelijk bedankt, Gerrit, for your inputs (including the idea of FIP) and for examining thoroughly the results with constructive criticism. You thought me to read the physics beyond the data and showed me positive thinking and great competence. For this thesis, I made large use of X-rays. In my 22 (!) synchrotron (ESRF) experiments, I benefitted of the help of many people. I am grateful to the staff of the beamlines BM26/ DUBBLE, ID02 and ID11 of the ESRF for supporting our experiments. Here, I would like to mention the persons that introduced me to X-ray scattering: Dr. Guido Heunen and Dr. Ann Terry. Guido cleverly arranged setups at BM26 at all times of the day and of the night. Ann is a tireless coworker that solved many setup and software deadlocks at ID02 and ID11. The discussions with Wim Bras were important in building motivations. I also acknowledge Dr. Giuseppe Portale for the support at BM26, Dr. Caroline Curfs for ID11 and of Dr. Peter Boesecke for ID02. Colleagues from SKT and M@te gave their valuable contribution to several ESRF experiments. Above all, I am indebted with Jan-Willem Housmans for all the time spent at the ESRF, including the always present dropjes. I had good time working in SKT where I benefitted of the support of several people. I would like to thank Dr. Han Goossens for the discussions and the comments, especially in the early stages of my PhD and Dr. Juan Fran Vega for teaching me the tricks of rheometry. My gratitude extends to Dirk Lippits for the stimulating discussions about rheology and for 133 calling me ‘maestro’, to Jules Harings for thorough discussions about my and his research that helped in building motivations, to Esther Vinken for the help with TGA and ‘that’ sms, to Nilesh Kukalyekar for the fruitful collaboration, to Irina Cotiuga for the discussions and the good time, to Joachim Loos, Anne Spoelstra and Pauline Schmit for microscopy, to Denka Hristova for the help in solving my initial problems with X-rays, to Sachin Jain for the discussions, to Joost Valeton for DSC, to Jules Kierkels for the help with X-rays and to Sainath Vaidya for the CNTs solutions. The list includes also my officemates at STO 0.42 Mano Prusty, Dirk Lippits and Gizela Mikova and the other people of SKT and PTG: Laurent Nelissen, Peter Koets, Mark van der Mee, Chunxia Sun, Lijing Xue, Rafiq Ahmed, Marjolein Diepens, Marjoleine Drieskens, Roy l’Abee, Saeid Talebi, Weizhen Li, Maya Ziari, Bjorn Teurlings, Edgar Karssenberg, Soney Varghese, Martijn Jansen, Cees Weijers, Thierry Leblanc. Thanks to Bob Fifield for all the incursions in my office, the handshakes and the always open debate about the reason why the Romans did not invade Wales. Thanks to Elly Langstad and Ineke Kollenburg for all the care. From the group M@te, I am grateful to Reinhard Forstner, Roel Janssen, Frederico Custodio and Rudi Steenbakkers for the discussions and the good time together, also at the ESRF. I also want to thank Sjef Garenfeld for providing help to my last minute requests. My life in Eindhoven was always hectic thanks to some memorable characters. I was introduced en la movida de Eindhoven by Carlos. Graciaaaas, Carlos, for being a real friend. All the events (including the FORT) together with Blanca, David, Ben and Chiara have something to be remembered. I enjoyed these times like there is no tomorrow. Unforgettable. This list cannot be closed without acknowledging the friendship of Dillip (is he spanish?), the class of Brett, the black clothing of Chris, the bbqs of Dani, the world according to Kuba, the dinners of Lorenzo, the sweetness of Mari, the twisted reasoning of Vincente, the French accent of Emilie, the smiling eyes of Amparo, the sincerity of Víctor and the smiles of Monica. Last and certainly not least, Chiara. I take the risk of being banal saying that you gave me incredible support. We had great time and you also tolerated my difficult moments. I thank you for the comprehension and the patience. Per chiudere, ringrazio la mia famiglia per tutto il supporto e gli incoraggiamenti ricevuti e quella Rabotti per la stima e le vacanze in Toscana. 134 Curriculum Vitae  The author was born in Pompei, Italy, on June 25th 1977. After finishing secondary school (Liceo Scientifico ‘E.Pascal’, Pompei), he studied Chemical Engineering at the University ‘Federico II’ of Naples (Italy). He completed his master thesis, entitled ‘Temperature effects in flow induced crystallization of thermoplastic polymers’ under the supervision of Prof. Nino Grizzuti. At the beginning of 2004, he started his Ph.D. in the Polymer Technology group (SKT) of Prof. P.J. Lemstra at Eindhoven University of Technology under the guidance of Prof. dr. S.Rastogi and dr. G.W.M. Peters. 135